Ductile Ultra High Strength Medium Manganese Steel Produced Through Continuous Annealing and Hot Stamping

- Colorado School of Mines

The present invention relates to the use of medium manganese steels for hot stamping to produce ultra high strength steel blanks with improved residual ductility of hot formed steels (as compared to boron-added steels such as 22MnB5 grade). The medium manganese steel sheets require much lower processing temperatures (<700° C. vs. >900° C.) prior to forming operation, allowing reduction of energy consumption during manufacturing.

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Description
CROSS REFERENCE TO RELATED APPLICATIONS

This application claims priority under 35 U.S.C. §119(e) to U.S. Provisional application Ser. No. 62/151,147, filed Apr. 22, 2015, which is incorporated herein in its entirety by reference.

FIELD OF THE INVENTION

The present invention relates to ductile ultra high strength medium manganese steels and methods of producing the same.

BACKGROUND

With worldwide pressure mounting to improve vehicle fuel economy and reduce carbon emissions to address climate change, improving the properties of high strength automotive steel remains an important task. Aggressive fuel economy legislation has been enacted for vehicles in the near future, and the U.S. government has recognized the importance of advanced high strength steels (AHSS). Property targets are still being developed for next generation AHSS, but in general there remains a focus on increasing the combination of tensile strength and tensile ductility, recognizing that “local formability” (hole expansion or sheared edge stretchability, shear fracture behavior or die radii, bendability etc.) also represents a key attribute that must be sufficient if implementation success is to be achieved.

Medium manganese transformation induced plasticity (TRIP) steels are promising 3rd generation AHSS with potential to exhibit excellent strength-ductility combinations relevant for automotive cold-forming applications. In contrast, low-alloy B-added steels (e.g. 22MnB5) are typically used for hot forming of sheet steels intended for use in structural components of auto bodies. These hot formed sheet steels exhibit ultra high strength in the final product. However, the residual ductility of the steels is quite low due to the martensitic microstructure, and higher ductility may be of interest to increase structural integrity during a crash.

Efforts have been made to increase the residual ductility of hot formed steels either by modifying the conventional hot forming thermal cycle (for B-free low-C steels) or by using a post-hot forming tempering treatment (for B-added steels). In the former case, more complex thermal cycle changes in hot forming may be necessary, and for the latter additional heat treating steps need to be added to existing forming lines. In either case, the improved elongation levels were reported to be limited to about 15%. Therefore, the development of an alternative class of hot formed steels that can fulfill the strength criterion required for automotive structural components and simultaneously possess enhanced elongation levels with no substantial changes in hot forming process is desired.

Press hardening, also called press forming, hot forming, or hot stamping, is a technique used to both form and harden a sheet into a final component. In this process, a blank sheet is heated to austenitizing temperatures and then simultaneously pressed and quenched in the dies. This serves to shape and harden the part by transforming the microstructure from austenite to martensite. This method is often used to form automotive parts which have complex shapes and require ultra high strength. Hot-stamping of uncoated sheets, however, leads to oxidation, therefore requiring oxide removal from the surface. Aluminum-silicon coatings have been used to address this problem, but Zn-based coatings may be preferable as they provide galvanic corrosion resistance. Zn-based coatings have potential, but the low melting temperatures of Zn and Zn—Fe compounds can lead to surface cracking during high temperature forming, associated with grain boundary penetration of liquid Zn. These cracks could compromise aesthetic standards and, more importantly, the integrity of the part. Currently, the mitigation of this cracking has focused on modifying the Zn coatings, rather than modifying the substrate steel.

SUMMARY

The present invention is directed towards a method of producing ductile, ultra high strength medium manganese steels produced through continuous annealing and hot stamping which has ultra high strength qualities and enhanced elongation levels. For structural components in automobiles such as B-pillars, ultra high strength is required to maintain a good structural integrity and crash performance, and simultaneously the component must be formed from steel sheets. Usually, an inverse relationship is observed between formability and strength of sheet metals. Currently available cold-forming technology can be used comfortably for steel sheets of strength up to 980 MPa, although some recent efforts suggest its successful application to sheets of strength of about 1200 MPa or higher.

Medium manganese TRIP steels contain manganese levels in the range of 2.5-15 mass % and may exhibit exceptional ductility (>30%) with reasonably high strength (600-1200 MPa). The carbon contents of these steels are normally low (<0.2 mass %), although stabilization of austenite is caused by both manganese and carbon. These steels are usually subjected to an intercritical annealing treatment to retain very high amounts of austenite (up to 60% by volume) at room temperature, without resorting to any additional heat treatment steps. Primarily, manganese partitions into austenite during the intercritical heat treatment ensuring the high amount and stability of austenite at room temperature. The final microstructure of medium Manganese TRIP steels contains a high amount of austenite in a ferritic matrix (duplex), occasionally with small amounts of tetragonal α′- or hexagonal ε-martensite. The austenite at room temperature is metastable and transforms to body-centered tetragonal α′-martensite upon application of load through displacive transformation causing a TRIP effect. The TRIP effect leads to high strain or work hardening rates, a delay in the onset of instability, and simultaneous increases in strength and ductility. Thus, the excellent mechanical properties of medium manganese TRIP steels make them attractive throughout industry.

Hot forming/stamping/pressing, whereby the sheets are formed at high temperatures and simultaneously press-quenched to room temperature, is used to produce the ultra high strength auto components. Subtle addition of boron is done to enhance the hardenability of hot forming steels so that a predominantly martensitic microstructure is formed in hot formed condition that ensures the ultra high strength. However, due to a martensitic microstructure, the residual ductility in hot formed components is very low (total elongation <5%). The present invention overcomes this drawback by providing a hot formed sheet steel product with both ultra high strength and high residual ductility.

Another benefit of the present invention is an increase in hardenability. By increasing manganese levels to reduce the austenitizing temperature, hardenability will also be increased substantially, thereby reducing the sensitivity of the hot-stamping process to cooling rate after forming. Titanium and boron additions may thus be no longer required, and if hardenability is increased sufficiently, then press quenching may not be required in order to achieve a martensitic final structure. If press-quenching could be avoided, the hot-stamping process could be conducted at much increased productivity and reduced capital costs.

The increase in manganese levels may also help broaden the use of AHSS in other areas such as automotive seating applications. For these applications, the sheet thickness is typically less than about 0.8 mm, and current grades are reportedly not rigid enough to be successfully transferred from the austenitizing furnace to the pressing dies. The inability to maintain the austenitizing temperature during transfer also results in low hardening during forming. Higher manganese levels would improve the high temperature strength of the sheet and make the transfer easier. Also, if a fully martensitic structure can be achieved at a slower cooling rate, maintaining elevated temperature during the transfer to the dies would have a less pronounced impact on hardening during press quenching.

The typical manganese levels in hot stamping alloys are on the order of 1.2-2 mass %, and provide sufficiently hardenability to achieve martensitic microstructures when die quenched at a sufficient rate (greater than about 30° C./s). In the present invention, the manganese levels are between about 2.5 mass % to about 15 mass %, which is greater than the typical amount of manganese found in alloy subjected to hot stamping. The levels of manganese in the present invention are employed to stabilize austenite, which thereafter retains austenite in the product at room temperature and enhances ductility in the final microstructure. The minimum amount of manganese to result in the austenite can be dependent upon the alloy used with the invention and the processing parameters. Further, the present invention can also lead to improved behavior of Zn-based coatings for use during hot forming. The hot forming of uncoated steel sheets leads to oxidation, requiring oxide removal from the surface. While aluminum-silicon coatings have been used to address this problem, Zn-coatings may be preferable as they provide galvanic corrosion resistance. However, because of the low melting temperatures of Zn and Zn—Fe compounds, surface cracking during high temperature forming can occur, associated with grain boundary penetration of liquid Zn. However, by modifying the substrate steel by enhancing the amount of austenite stabilizing elements, in this case manganese, the processing temperatures can be significantly reduced and liquation of the coating can be avoided and cracking may be eliminated.

Utilizing the method of the present invention, it is possible to produce a ductile, ultra high strength 10Mn steel with high retained austenite content by continuous annealing the cold rolled steel at various temperatures. While not wanting to be bound by theory, the transformation induced plasticity effect due to the presence of metastable austenite along with strengthening from an ultrafine grain size and high dislocation density due to short non-equilibrium processing are believed to contribute to the ultra high strength achieved in the 10Mn steel. Moreover, subjecting the 10Mn TRIP steel to hot forming thermal cycles at various temperatures after different continuous annealing treatments yields ultra high strength levels with higher residual total elongation than traditional B-added hot formed steels.

An aspect of the present invention is a method for producing ultra high strength steel. The method includes the steps of providing a ferrous alloy comprising carbon and manganese, annealing the ferrous alloy at a first temperature to form an annealed alloy, hot forming the annealed alloy at a second temperature to form an intercritical or austenitic microstructure, and cooling the austenitic alloy to room temperature to form an ultra high strength steel.

An aspect of the present invention is an ultra high strength steel product. The steel product has a composition of between about 2.5 mass % and 15 mass % manganese, and wherein the matrix microstructure comprises about 5 vol. % to about 65 vol. % austenite, about 35 vol. % to about 90 vol. % ferrite, and about 0 vol. % to about 95 vol. % martensite.

An aspect of the present invention is a method for producing an ultra high strength steel. The method includes the steps of providing a hot-rolled medium manganese steel comprising between about 2.5 mass % and about 15 mass % manganese to produce a hot rolled steel, hot band annealing the hot rolled steel in the temperature range of between about 600° C. and about 720° C., cold-rolling the material to reduce the thickness by between about 30% and about 70%, continuously annealing the cold-rolled material at a temperature between about 600° C. and about 950° C. to form an annealed steel, hot forming the annealed steel at a temperature between about 600° C. and about 950° C. to form a hot formed steel, and quenching the hot formed steel to form the ultra high strength steel.

An aspect of the invention is a method for producing ultra high strength steel alloy. The method includes providing a ferrous alloy comprising carbon and greater than 2.5 mass manganese, annealing the ferrous alloy at a temperature to form an annealed alloy comprising an austenitic microstructure, and cooling the steel comprising the austenitic microstructure to room temperature to form an ultra high strength steel comprising the austenitic microstructure, and at least one microstructure of ferrite or martensite.

An aspect of the invention is a method for producing ultra high strength steel. The method includes providing a cold-rolled medium manganese steel comprising between about 2.5 mass % and about 12 mass % manganese, hot forming the steel at a temperature between about 600° C. and about 900° C. to form a hot formed steel, and quenching the hot formed steel to form the ultra high strength steel.

These and other advantages will be apparent to one skilled in the art in view of this disclosure.

BRIEF DESCRIPTION OF THE FIGURES

FIG. 1 is a flowchart displaying the pre-processing treatment of the as-received cold-rolled medium manganese sheets;

FIG. 2 depicts an engineering stress-strain curve of the continuously annealed cold- rolled medium manganese sheets;

FIG. 3(a) depicts a graph illustrating the ultimate tensile strength of the continuously annealed samples plotted as a function of annealing temperature

FIG. 3(b) depicts a graph illustrating the total elongation of the continuously annealed samples plotted as a function of annealing temperature;

FIG. 4(a) depicts engineering stress-strain curves of the hot formed samples for reheating at 650° C.;

FIG. 4(b) depicts engineering stress-strain curves of the hot formed samples for reheating at 700° C.;

FIG. 5(a) illustrates a SEM image illustrating the typical microstructures for continuously annealed (CA) (700° C. CA);

FIG. 5(b) illustrates a SEM image illustrating the typical microstructures for hot formed (HF) (700° C. CA+700° C. HF) conditions;

FIG. 6(a) depicts a graph of the austenite fractions of the continuously annealed measured by x-ray diffraction (XRD);

FIG. 6(b) depicts a graph of the austenite fractions of the hot formed samples measured by XRD;

FIG. 7(a) depicts a graph illustrating the average full widths at half maxima (FWHM) of all the ferrite peaks at various temperatures from the XRD patterns for the continuously annealed conditions; and

FIG. 7(b) depicts a graph illustrating the average FWHM of all the ferrite peaks at various hot forming conditions from the XRD patterns.

 DETAILED DESCRIPTION

The present invention relates to a novel method of producing ultra high strength, medium manganese steel through hot forming. In one embodiment, the medium manganese steel is subjected to heat treatments of a typical direct hot forming route in muffle furnaces.

An aspect of the present invention is a method for producing ultra high strength steel. The method including the steps of providing a ferrous alloy comprising carbon and manganese, annealing the ferrous alloy at a first temperature to form an annealed alloy, hot forming the annealed alloy at a second temperature to form a twice annealed alloy, and cooling the twice annealed alloy to room temperature to form an ultra high strength steel comprising austenite in the microstructure.

In some embodiments of the present invention, annealing of the ferrous alloys can be performed at a temperature of between about 600° C. and about 950° C. Other suitable temperature ranges can include between 650° C. and 800° C., between 600° C. and about 750° C., between about 750° C. and about 900° C., or between 750° C. and 800° C. The annealing can be a continuous anneal. The intercritical or austenitic annealing temperature can be dependent upon the particular alloy. By way of non-limiting example only, the intercritical annealing temperature of a 10Mn alloy can be between about 600° C. to about 750° C., and the austenitic annealing temperature range of between about 750° C. to about 800° C., or above. Throughout the

Specification, the intercritical or austenitic annealing temperature ranges of 10Mn alloy will be discussed, though one skilled in the art would understand that suitable intercritical and austenitic annealing temperature ranges can be determined by one skilled in the art without deviating from the present invention. The annealing can utilize a heating rate of about 7° C. per second or in case of furnace heating or faster in case of induction heating of the alloy (greater than about 100° C./s). The heating rate can slightly influence manganese partitioning between the phases, and perhaps the scale of the microstructure in the heating temperature, thereby influencing the final microstructure and properties. At least one annealing cycle can be used, though it is understood that multiple annealing cycles could be performed without deviating from the invention. The soaking time of the annealing can be between about 60 seconds and about 10 minutes, in some embodiments about 3 minutes. Excessive annealing times would not be expected to diminish the applicability in the intercritical regime with respect to properties, but would be expected to reduce productivity, or influence the behavior of metallic coatings, etc.

An important advantage of the present invention is that the resulting alloy material resulting from the method of the present invention can produce an alloy comprising high amount of austenite without requiring a high cooling rate of about 30° C./s or faster for a traditional hot stamping alloy (e.g. 22MnB5). Rather, the manganese content of the alloy allows the required alloy to be formed at cooling rates that can be less than about 30° C./s, in some embodiments between about 0.03° C./s and about 5° C./s. Because the cooling rate is greatly reduced with the present invention, expensive equipment required to, for example, extract heat from the dies used in the hot forming/hot stamping process and high energy consumption can be reduced. Furthermore, the distortion or spring back can be reduced with modified tooling or fixturing during the hot stamping process.

In some embodiments, the alloy can be hot formed following the annealing. The annealed alloy can be cooled to room temperature prior to the hot forming the annealed alloy, though one skilled in the art would understand that the material need not be cooled before hot forming the alloy. The hot forming can be performed at a temperature of between about 600° C. and about 950° C., preferably where the hot forming temperature can be at the intercritical annealing temperature range (e.g. between about 600° C. and 750° C.). The hot forming temperature can also be at the austenitizing annealing temperature (e.g. between about 750° C. and about 950° C.). The material can be held at temperature for between about 1 minute and about 10 minutes, in some embodiments up to about 8 minutes, in some embodiments about 3 minutes.

In some embodiments, the total mass of the ferrous alloy can comprise between about 2.5 mass % and about 15 mass %, between about 2.5 mass % and about 12 mass % manganese, between about 3 mass % and 10 mass % manganese, between about 3 mass % and about 12 mass % manganese, between about 7 mass % and about 10 mass %, between about 8 mass % and about 10 mass % manganese, or between about 8 mass % to about 12 mass % manganese. This amount of manganese can allow for a product that has high tensile strength (by the production of martensite) and high ductility (by the production of austenite).

The ultra high strength steel comprises at least one microstructure of austenite, martensite, ferrite, or combinations thereof. The formation of martensite or ferrite can depend on the annealing temperature. For example, if the steel is formed over a temperature from about 600° C. to about 750° C. (intercritical annealing), then the microstructure of austenite, martensite and ferrite can be formed. If the annealing temperature is formed over a temperature between about 750° C. and about 800° C. (austenitic annealing), then the steel can comprise microstructures of austenite and martensite. The microstructures formed and the applicable processing temperatures can depend on the exact alloy used with the invention. Austenite can be present in the ultra high strength steel at a volume percentage greater than about 10%, in some embodiments between about 18.5 vol. % and about 56 vol. %, with the remainder being ferrite and/or martensite. One skilled in the art would understand that the amount of austenite in the sample would affect the properties of the ultra high strength steel. One skilled in the art would also understand that the amount of austenite present in the final product would be dependent upon processing parameters as discussed in more detail herein. The stability of the austenite can also affect the properties of the ultra high strength steel. These stability properties can be both composition and process dependent.

At least one additive can be added to the steel material prior to annealing. Suitable additives include, but are not limited to, aluminum, silicon, carbon, microalloying elements such as titanium, niobium, or vanadium and combinations thereof. Silicon levels (up to about 3 mass %, in some embodiments between about 0.1 mass % and about 1.5 mass % silicon) can increase strength and potentially modify iron-carbide precipitation kinetics. Aluminum levels (up to about 3 mass %, or potentially more, in some embodiments between about 0.5 mass % and about 3 mass % aluminum) can modify the phase stability and intercritical annealing response as well as to beneficially influence the hydrogen resistance of the austenitic phase. Microalloying with niobium, vanadium or titanium (up to levels of about 0.2 mass %) can refine the microstructure and increase the strength through alloy carbonitride precipitation. Carbon variations up to approximately 0.4 mass %, in some embodiments between 0.05 mass % and about 0.3 mass % carbon, can increase the strength of the martensite, and can achieve substantially higher strength for some applications that can benefit from higher strengths. Titanium, nitrogen, and boron levels can be modified to influence solute nitrogen levels and grain boundary cohesion, as well as to influence potential hydrogen trapping behavior. The balance of the alloy can be iron with inevitable impurities.

Some embodiments of the invention can further include a hot band annealing process. The hot band annealing can occur following the coiling step after hot rolling. The hot band annealing can occur at a temperature between about 600 ° C. and about 720° C., in some embodiments about 640° C. for between about 60 minutes and about 168 hours, in some embodiments about 96 hours. The hot band annealing can occur in a hydrogen gas atmosphere. Following the hot band annealing, the steel can be air cooled to room temperature.

In some embodiments of the present invention, the ultra high strength steel product can result in at least one property. For example, the product can result in a grain size of between about 0.1 micron and about 5 micron. In some embodiments, a dislocation density property of the steel can be between about 109 m−2 and about 1010 m−2. In some embodiments, the tensile strength property of the steel product can be between about 1000 MPa and about 1700 MPa. The residual total elongation property of the steel can be between about 10% and about 30%. In other embodiments, the x-ray diffraction pattern of the product can have an average full widths at half maxima (FWHM) of all ferrite peaks of between about 0.25° and about 1°. Additional property ranges can be accessed using the present invention by additional alloy modifications such as variations in the carbon concentration, annealing temperatures, etc.

An aspect of the present invention is an ultra high strength steel product. The composition of the product comprises between about 2.5 mass % and about 15 mass % manganese, and the balance being iron and inevitable impurities. The product results in a microstructure that comprises from about 5 vol. % to about 65 vol. % austenite, and the balance martensite. In some embodiments, the product can comprise between about 35 vol. % to 90 vol. % ferrite.

In some embodiments, retained austenite levels above about 5 vol. % can provide enhanced work hardening due to the TRIP or other effects, thereby enhancing the ductility at high strength levels. Austenite can be present in the ultra high strength steel at a volume percentage greater than about 10%, in some embodiments between about 18.5 vol. % and about 56 vol. %, with the remainder being ferrite and/or martensite. One skilled in the art would understand that the amount of austenite in the sample would affect the properties of the ultra high strength steel. One skilled in the art would also understand that the amount of austenite present in the final product would be dependent upon processing parameters as discussed in more detail herein. The stability of the austenite can also affect the properties of the ultra high strength steel. These stability properties can be both composition and process dependent.

In some embodiments, the total mass of the ferrous alloy can comprise between about 2.5 mass % and about 15 mass %, between about 2.5 mass % and about 12 mass % manganese, between about 3 mass % and 10 mass % manganese, between about 3 mass % and about 12 mass % manganese, between about 7 mass % and about 10 mass %, between about 8 mass % and about 10 mass % manganese, or between about 8 mass % to about 12 mass % manganese. This amount of manganese can allow for a product that has high tensile strength (by the production of martensite) and high ductility (by the production of austenite).

In some embodiments of the present invention, the ultra high strength steel product can result in at least one property. For example, the product can result in a grain size of between about 0.1 micron and about 5 micron. In some embodiments, a dislocation density property of the steel can be between about 109 m−2 and about 1010 m−2. In some embodiments, the tensile strength property of the steel product can be between about 1000 MPa and about 1700 MPa. The residual total elongation property of the steel can be between about 10% and about 30%. In other embodiments, the x-ray diffraction pattern of the product can have an average FWHM of all ferrite peaks of between about 0.25° and about 1°. Additional property ranges can be accessed using the present invention by additional alloy modifications such as variations in the carbon concentration, annealing temperatures, etc.

At least one additive can be added to the steel material prior to annealing. Suitable additives include, but are not limited to, aluminum, silicon, carbon, microalloying elements such as titanium, niobium, or vanadium and combinations thereof. Silicon levels (up to about 3 mass %, in some embodiments between about 0.1 mass % and about 1.5 mass % silicon) can increase strength and potentially modify iron-carbide precipitation kinetics. Aluminum levels (up to about 3 mass %, or potentially more, in some embodiments between about 0.5 mass % and about 3 mass % aluminum) can modify the phase stability and intercritical annealing response as well as beneficially influence the hydrogen resistance of the austenitic phase. Microalloying with niobium, vanadium or titanium (up to levels of about 0.2 mass %) can refine the microstructure and increase the strength through alloy carbonitride precipitation. Carbon variations up to approximately 0.4 mass %, in some embodiments between 0.05 mass % and about 0.3 mass % carbon, can increase the strength of the martensite, and can achieve substantially higher strength for some applications that can benefit from higher strengths. Titanium, nitrogen, and boron levels can be modified to influence solute nitrogen levels and grain boundary cohesion, as well as to influence potential hydrogen trapping behavior. The balance of the alloy can be iron with inevitable impurities.

An aspect of the present invention is a method for producing ultra high strength steel. The method comprises the steps of providing a cold-rolled medium manganese steel comprising between about 2.5 mass % and about 15 mass % manganese, continuously annealing the cold-rolled medium manganese steel at a temperature between about 600° C. and about 950° C. to form an annealed steel, hot forming the annealed steel at a temperature between about 600° C. and about 950° C. to form a hot formed steel, and quenching the hot formed steel to form the ultra high strength steel.

In some embodiments of the present invention, annealing of the ferrous alloys can be performed at a temperature of between about 600° C. and about 950° C. Other suitable temperature ranges can include between 650° C. and 800° C., between 600° C. and about 750° C., between about 750° C. and about 900° C., or between 750° C. and 800° C. The intercritical or austenitic annealing temperature can be dependent upon the particular alloy. By way of non-limiting example only, the intercritical annealing temperature of a 10Mn alloy can be between about 600° C. to about 750° C., and the austenitic annealing temperature range of between about 750° C. to about 800° C. Throughout the Specification, the intercritical or austenitic annealing temperature ranges of 10Mn alloy will be discussed, though one skilled in the art would understand that suitable intercritical and austenitic annealing temperature ranges can be determined by one skilled in the art without deviating from the present invention. The annealing can utilize a heating rate of about 7° C. per second in case of furnace heating or faster in case of induction heating of the alloy (greater than about 100° C./s). The heating rate can slightly influence manganese partitioning between the phases, and perhaps the scale of the microstructure in the heating temperature, thereby influencing the final microstructure and properties. At least one annealing cycle can be used, though it is understood that multiple annealing cycles could be performed without deviating from the invention. The soaking time of the intercritical annealing can be between about 60 seconds and about 10 minutes, in some embodiments about 3 minutes. Excessive annealing times would not be expected to diminish the applicability in the intercritical regime with respect to properties, but would be expected to reduce productivity, or influence the behavior of metallic coatings, etc.

An important advantage of the present invention is that the resulting alloy material resulting from the method of the present invention can produce an alloy comprising high amount of austenite without requiring a high cooling rate of about 30° C./s or faster for a traditional hot stamping alloy (e.g. 22MnB5). Rather, the manganese content of the alloy allows the required alloy to be formed at cooling rates that can be less than about 30° C./s, in some embodiments between about 0.03° C./s and about 5° C./s. Because the cooling rate is greatly reduced with the present invention, expensive equipment required to, for example, extract heat from the dies used in the hot forming/hot stamping process and high energy consumption can be reduced. Furthermore, the distortion or spring back can be reduced with modified tooling or fixturing during the hot stamping process.

In some embodiments, the alloy can be hot formed following the annealing. The annealed alloy can be cooled to room temperature prior to the hot forming the annealed alloy, though one skilled in the art would understand that the material need not be cooled before hot forming the alloy. The hot forming can be performed at a temperature of between about 600° C. and about 950° C., preferably where the hot forming temperature can be at the intercritical annealing temperature range (e.g. between about 600° C. and 750° C.). The hot forming temperature can also be at the austenitizing annealing temperature (e.g. between about 750° C. and about 950° C.). The material can be held at temperature for between about 1 minute and about 10 minutes, in some embodiments up to about 8 minutes, in some embodiments about 3 minutes.

In some embodiments, the total mass of the ferrous alloy can comprise between about 2.5 mass % and about 15 mass %, between about 2.5 mass % and about 12 mass % manganese, between about 3 mass % and 10 mass % manganese, between about 3 mass % and about 12 mass % manganese, between about 7 mass % and about 10 mass %, between about 8 mass % and about 10 mass % manganese, or between about 8 mass % to about 12 mass % manganese. This amount of manganese can allow for a product that has high tensile strength (by the production of martensite) and high ductility (by the production of austenite).

The ultra high strength steel comprises at least one microstructure of austenite, martensite, ferrite, or combinations thereof. The formation of martensite or ferrite can depend on the annealing temperature. For example, if the steel is formed over a temperature from about 600° C. to about 750° C. (intercritical annealing), then the microstructure of austenite, martensite and ferrite can be formed. If the annealing temperature is formed over a temperature between about 750° C. and about 800° C. (austenitic annealing), then the steel can comprise microstructures of austenite and martensite. The microstructures formed and the applicable processing temperatures can depend on the exact alloy used with the invention. Austenite can be present in the ultra high strength steel at a volume percentage greater than about 10%, in some embodiments between about 18.5 vol. % and about 56 vol. %, with the remainder being ferrite and/or martensite. One skilled in the art would understand that the amount of austenite in the sample would affect the properties of the ultra high strength steel. One skilled in the art would also understand that the amount of austenite present in the final product would be dependent upon processing parameters as discussed in more detail herein. The stability of the austenite can also affect the properties of the ultra high strength steel. These stability properties can be both composition and process dependent.

Additives can be added to the steel material prior to annealing. Suitable additives include, but are not limited to, aluminum, silicon, carbon, microalloying elements such as titanium, niobium, or vanadium and combinations thereof. Silicon levels (up to about 3 mass %, in some embodiments between about 0.1 mass % and about 1.5 mass % silicon) can increase strength and potentially modify iron-carbide precipitation kinetics. Aluminum levels (up to about 3 mass %, or potentially more, in some embodiments between about 0.5 mass % and about 3 mass % aluminum) can modify the phase stability and intercritical annealing response as well as to beneficially influence the hydrogen resistance of the austenitic phase. Microalloying with niobium, vanadium or titanium (up to levels of about 0.2 mass %) can refine the microstructure and increase the strength through alloy carbonitride precipitation. Carbon variations up to approximately 0.4 mass %, in some embodiments between 0.05 mass % and about 0.3 mass % carbon, can increase the strength of the martensite, and can achieve substantially higher strength for some applications that can benefit from higher strengths. Titanium, nitrogen, and boron levels can be modified to influence solute nitrogen levels and grain boundary cohesion, as well as to influence potential hydrogen trapping behavior. The balance of the alloy can be iron with inevitable impurities.

Some embodiments of the invention can further include a hot band annealing process. The hot band annealing can occur following the coiling step after hot rolling. The hot band annealing can occur at a temperature between about 600° C. and about 720° C., in some embodiments about 640° C. for between about 60 minutes and about 168 hours, in some embodiments about 96 hours. The hot band annealing can occur in a hydrogen gas atmosphere. Following the hot band annealing, the steel can be air cooled to room temperature.

In some embodiments of the present invention, the ultra high strength steel product can result in at least one property. For example, the product can result in a grain size of between about 0.1 micron and about 5 micron. In some embodiments, a dislocation density property of the steel can be between about 109 m−2 and about 1010 m−2. In some embodiments, the tensile strength property of the steel product can be between about 1000 MPa and about 1700 MPa. The residual total elongation property of the steel can be between about 10% and about 30%. In other embodiments, the x-ray diffraction pattern of the product can have an average FWHM of all ferrite peaks of between about 0.25° and about 1°. Additional property ranges can be accessed using the present invention by additional alloy modifications such as variations in the carbon concentration, annealing temperatures, etc.

In general, the amount of metastable austenite in microstructures at room temperature can be a key factor in determining the mechanical properties of AHSS . Therefore, the evolution of tensile strength and elongation after continuous annealing of the material can be interpreted in terms of the retained austenite in the microstructures. With an increase in annealing temperature in the intercritical phase field of the 10Mn steel, the austenite present at the annealing temperature typically increases, while the stability of intercritical austenite typically decreases due to reduced equilibrium partitioning of manganese in the austenite.

In addition, through short continuous annealing of the relatively higher manganese-containing steel, a high amount of metastable austenite can be obtained in the microstructure. The tensile properties achieved following continuous annealing are unexpectedly high. While not wanting to be bound by theory, it is believed that the tensile properties may be high due to ultrafine grain size, and a higher dislocation density due to short annealing of the cold rolled material in addition to the presence of the metastable austenite in sufficient amounts in the microstructure.

Further, once the steels have been hot formed, the austenite content at room temperature in the microstructure decreases with increasing reheat temperature. The austenite contents indicate that probably the effect of the second annealing during hot forming may be more dominant in controlling the extent of manganese partitioning at high temperature than the initial microstructure developed after the first continuous annealing. Similarly, the elongation of the steel after hot forming may likely substantially affected by the austenite contents present in the microstructure.

The temperatures used for the hot forming are also much lower due to suppression of Ae3 temperature of the steel by high amount of manganese. The hot forming temperature providing improved mechanical properties can be at least about 200° C. lower than the temperature used for austenitizing 22MnB5, which is typically greater than about 900° C. This lower hot forming temperature allows for significant energy saving, reducing the already high cost associated with hot forming operation, and may have a favorable influence on the behavior of a galvanized coating.

Moreover, the increased levels of manganese in the steel reduce the austenitizing temperature, and leads to a substantial increase in hardenability, also reducing the sensitivity of the hot-stamping process to cooling rate after forming. Titanium and boron additions may thus be no longer required, and if hardenability is increased sufficiently, then press quenching may not be required in order to achieve a martensitic final microstructure. If press-quenching could be avoided, the hot-stamping process could be conducted at much increased productivity and reduced capital cost. Distortion may become more important if press quenching is eliminated, but this issue might be addressed easily by fixturing. That is, slower cooling in a simple fixture that restrains distortion can be accomplished, avoiding the more costly process of press-quenching where the stamping dies are typically engineered with internal cooling.

All publications, patents, and patent documents cited herein are incorporated by reference herein, as though individually incorporated by reference. The invention has been described with reference to various specific and preferred embodiments and techniques. However, it should be understood that many variations and modifications may be made while remaining within the spirit and scope of the invention.

It is appreciated that certain features of the invention, which are, for clarity, described in the context of separate embodiments, may also be provided in combination in a single embodiment. Conversely, various features of the invention that are, for brevity, described in the context of a single embodiment, may also be provided separately or in any sub-combination.

This invention now being generally described will be more readily understood by reference to the following examples, which are included merely for the purpose of illustration of certain aspects of the embodiments of the invention. The examples are not intended to limit the invention, as one of skill in the art would recognize from the above teachings and the following examples that other techniques and methods can satisfy the claims and can be employed without departing from the scope of the claimed invention.

EXAMPLES

The following examples utilized medium manganese steel (identified as 10Mn steel) which contained 9.76Mn-0.16C-1.37A1-0.19Si-0.0025N-0.0018S (mass %), and was received as 1 mm-thick cold rolled (CR) sheets. In a preferred embodiment, the CR sheets were processed as illustrated in FIG. 1. The CR sheets then underwent a continuous annealing heat treatment followed by a hot forming heat treatment.

In a typical continuous annealing cycle, steel sheet is uncoiled and passes through a preheated furnace before being re-coiled after annealing. While passing through the hot furnace, the steel sheet is heated up rapidly and soaked at the intended annealing temperature for a short time (1-3 min), depending on the line speed at which the coil passes through the furnace. Therefore, as will be seen below a short annealing time was used to approximately simulate the continuous annealing cycle.

In a direct hot forming cycle, as-supplied annealed steel blanks are heated (conventionally in the austenitic temperature range) again a few minutes (between about 3 and 8 minutes) and then transferred to the adjacent press, where the blank is formed into parts and simultaneously press-quenched at a high rate (>30 ° C./s). Therefore, as described below, to simulate the hot forming thermal cycle, prior continuous annealed steel sheets were subjected to a short annealing treatment followed by rapid quenching. It is to be noted here that although the process has been referred to as hot forming in this study, only the thermal cycles associated with hot forming were simulated; no deformation was induced to simulate ‘forming’ which is assumed to be acceptable for the purpose of demonstrating the inventive concept.

The continuous annealing temperature (TCA) was varied to find the optimum temperature for obtaining the best as-annealed mechanical properties, and the hot forming temperature (THF) was also varied for a particular continuous annealing temperature to obtain the best properties after hot-stamping. The use of different prior continuous annealing temperatures for a particular hot forming cycle would thus provide knowledge of the effect of prior microstructure on the development of microstructure after hot forming.

EXAMPLE 1 Mechanical Properties and Microstructure of CA Samples and HF Samples

The CR sheets were first subjected to heat treatments of a typical direct hot forming route in muffle furnaces. The as-received sheets were continuously annealed at 650° C., 700° C., 750° C., and 800° C. with a heating rate of approximately 7° C./s and a soaking time of 3 minutes followed by water quenching. The continuously annealed samples will be referred to as “CA” samples.

The CR sheets were also subjected to continuously annealing as described above. The CA sheets at each temperature were then further subjected to hot forming thermal cycles by reheating at 650° C., 700° C., 750° C., or 800° C. using the same heating rate and soaking time as in the continuous annealing treatments. The reheated samples were then quenched in water to simulate press-quenching. The hot formed samples will be referred to as “HF” samples. Two temperatures (650° C. and 700° C.) were selected from the intercritical phase field, and the other two (750° C. and 800° C.) from the austenitic phase field of the 10MnAl steel based on previous thermodynamic calculations. While 750° C. is very close to the temperature at which a sample would be fully austenitic (Ae3), 800° C. lies sufficiently above (70° C.) Ae3 and therefore allows study of the effect of austenitic grain growth on the hardenability and resulting final microstructure of the alloy. Hardenability of this 10 mass % manganese-containing steel is expected to be very high.

The mechanical properties were measured by room temperature tensile tests conducted with an initial strain rate of about 4.2×10−3 s−1 on ASTM E8 sub-sized specimens using an electromechanical test frame. XRD patterns of the heat treated specimens were obtained using Cu-Kα radiation at a 2θ range of 35-105° (θ=angle of diffraction) with a scanning step size of 0.008° and time per step of 19 seconds. Rietveld analyses of the XRD patterns were used to estimate the retained austenite fractions in the samples following the SAE SP-453 standard of the Society of Automotive Engineers. The FWHM of the XRD peaks were also determined to provide an indication of strain in the material. To avoid any effect of potential decarburization and oxidation of the samples during heat treating, material from the surface of the XRD specimens was removed by metallographic grinding followed by chemical polishing for at least 10 minutes in a solution of distilled H2O, H2O2, and HF (50:50:1 by volume). The microstructures of the metallographic samples etched in 2 vol. % nital were observed on the plane of rolling and normal directions (RD-ND plane) in a field emission scanning electron microscope (FESEM) using an accelerating voltage of 5 kV and 89 μA probe current.

Results: Mechanical Properties

The mechanical properties derived from the tensile data of CA samples are summarized in Table 1 while the engineering stress-strain curves of continuously annealed specimens are illustrated in FIG. 2. FIG. 2 clearly illustrates that the sample annealed at 700° C. has the highest amount of work hardening among all the annealing conditions whereas the sample annealed at 650° C. has the lowest work hardening. The samples annealed at or above 750° C. fractured in a brittle manner. These observations are reflected in the mechanical properties given in Table 1.

With increasing CA temperature, both the ultimate tensile strength (UTS) and tensile elongation (TE) initially increased, exhibiting a peak at 700° C., before decreasing. The sample annealed at 700° C. exhibited the highest UTS of 1619 MPa with reasonably high TE of 16.5% due to its high work hardening rate. The product of UTS and TE, which is considered as the energy absorption capability of the material, was also highest in the sampled annealed at 700° C., while the lowest values were obtained after annealing at higher temperatures (>750° C.). The UTS and TE are plotted as a function of annealing temperature in FIGS. 3(a) and 3(b) to visualize the trend illustrated in Table 1.

TABLE 1 Annealing YS UTS UE TE UTS × TE Temperature (° C.) (MPa) (MPa) (%) (%) (MPa. %) 650 1163 1163 13.9 14.9 17329 700 777 1619 15.5 16.5 26714 750 850 1366 2.5 2.5 3415 800 930 1188 1.3 1.3 1544

Where YS is yield strength, UE is ultimate elongation, TE is tensile elongation and UTS is ultimate tensile strength.

The mechanical properties of all the example HF samples are given in Table 2, while representative engineering stress-strain curves of the HF samples for reheating at 650° C. and 700° C. are depicted in FIGS. 4(a) and 4(b). Irrespective of prior CA treatment, the samples reheated at higher temperatures (750° C. and 800° C.) exhibited only modest ductility with ultra high UTS, and the samples subjected to HF cycles at lower temperatures (650° C. and 700° C.) showed a combination of high UTS (1104-1448 MPa) and high ductility (16.7-44.3% TE). The work hardening rates in samples reheated at 650° C. and 700° C., depicted in FIGS. 4(a) and 4(b), were also very high. It is interesting to observe that the samples which showed very low levels of ductility after CA treatment (for annealing at >750° C.; shown in Table 1), recorded a gain in elongation after the HF cycles, particularly for reheating at 650° C. and 700° C. (shown in Table 2). The best combination of ultra high tensile strength (UTS) (>1330-1448 MPa) and high ductility (>16.7-25.3%) was achieved for reheating at 700° C. for all prior CA conditions. The value of UTS×TE decreased with increasing reheat temperature.

TABLE 2 Annealing Hot Forming Temperature Temperature YS UTS UE TE UTS × TE (° C.) (° C.) (MPa) (MPa) (%) (%) (MPa · %) 650 650 1059 1153 43.4 44.3 51078 700 793 1367 24.3 25.1 34312 750 954 1522 4.1 4.1 6240 800 929 1293 1.9 1.9 2457 700 650 779 1395 21.0 21.3 29714 700 585 1448 16.7 16.7 24182 750 835 1605 7.3 7.4 11877 800 831 1469 3.5 3.5 5142 750 650 824 1202 30.1 31.4 37743 700 682 1417 21.4 22.3 31600 750 636 1568 8.0 8.7 13642 800 749 1501 4.3 4.3 6454 800 650 800 1104 28.4 33.0 36432 700 490 1330 21.1 22.1 29393 750 687 1561 4.8 4.9 7649 800 866 1348 2.4 2.4 3235

Results: Microstructure

The typical microstructures for the CA (700° C. CA) and HF (700° C. CA+700° C. HF) conditions are shown in the SEM images in FIGS. 5(a) and 5(b). Ultrafine ferrite and martensite-austenite (MA) constituents (associated with austenite at intercritical temperatures) were observed in both conditions. The grain size was in the submicron range, although a slight increase in grain size was observed after application of the HF cycles. The austenite fractions of the CA and HF samples measured by XRD, illustrated in FIGS. 6(a) and 6(b) respectively, decreased with increasing treatment temperature.

The average FWHM of all the ferrite peaks ((110), (200), (211), and (220)) from the XRD patterns are shown for the CA and HF conditions in FIGS. 7(a) and 7(b), respectively. The FWHM of ferrite for a fully annealed sample (650° C./32h, furnace cooled), determined using the same method, is also included in these FIGS. 7(a) and 7(b) for comparison. The ferrite FWHM is used here as an indicator of peak broadening due to strain. It is expected that peak broadening due to particle size was absent at the applicable grain size (submicron size), and the instrumental contribution to broadening would be the same in all the samples. Therefore, the FWHM values can be used for qualitative indication of strain. The ferrite FWHM values of both sets of samples were higher than that of the fully annealed sample indicating higher strain the CA and HF samples, likely associated with incomplete recrystallization along with martensite formation during final cooling of intercritical austenite.

Analysis

The amount and stability of austenite, grain size, and dislocation density in the microstructures are key factors that determine the mechanical properties of the investigated 10Mn steel. The sub-micron ultrafine grain size in all the samples (illustrated in FIGS. 5(a) and 5(b)), was a result of the relatively short treatment times and the duplex microstructure. The UTS for both the CA and HF conditions of the 10Mn steel was unusually high, and, without wishing to be bound to any particular theory, may be related to the Hall-Petch hardening effect of the ultrafine grain size of the constituent phases in combination with higher dislocation density in the samples, along with important effects from austenite.

In medium manganese steels, with a decrease in annealing temperature a higher degree of manganese partitioning can occur in austenite, and therefore higher amounts of retained austenite were obtained in the samples continuously annealed at lower temperatures (FIG. 6(a)). The austenite stability decreases due to reduced manganese partitioning in austenite at higher temperatures. The sample annealed at 700° C. exhibited the highest work hardening rate among the CA samples leading to a good combination of UTS and TE (as shown in Table 1), probably due to an optimum metastability of austenite.

The two samples annealed at higher temperatures (750° C. and 800° C.) exhibited low ductility due a variety of reasons. The sample annealed at 750° C. possessed a relatively high amount of austenite (FIG. 6(a)) but manganese enrichment in austenite is expected to be low causing early transformation of austenite to martensite during deformation. By contrast, the sample annealed at an even higher temperature (800° C.) had a predominantly martensitic microstructure (with low ferrite content, and low measured austenite content, as shown in FIG. 6(a)), and thus exhibited low ductility.

The austenite fractions in the HF samples decreased with increasing reheat temperature (as illustrated in FIG. 6(b)) due to the reasons discussed above. The austenite contents, particularly the higher measured austenite fractions in the samples annealed at higher temperatures but subjected to lower reheat temperatures, indicate that the effect of the second annealing during reheating is probably more dominant in controlling the extent of manganese partitioning than the initial microstructure developed after continuous annealing.

The samples reheated at 750° C. exhibited a low amount of austenite (as shown in FIG. 6(b)) along with expected low austenite stability as discussed above. Therefore, the observed peak UTS in these samples, shown in Table 2, is possibly due to the presence of unstable austenite and martensite in the microstructure. The UTS decreased after reheating at 800° C. due to the absence of necking during tensile deformation as low austenite contents in FIG. 6(b) indicate a high amount of martensite in these samples. The elongation values of the HF samples appear to substantially correlate with their austenite fractions.

The mechanical properties achieved after HF cycles in the 10Mn steel are particularly interesting. The results shown in Table 2 clearly suggest that it is possible to achieve UTS and TE in the ranges of 1330-1448 MPa and 16.7-25.3% (mostly >22.1% except for one HF condition) respectively after applying HF cycles at 700° C. Comparing these mechanical properties to the typical values obtained in the currently used hot forming steel grade 22MnB5 (UTS>1500 MPa, TE=5%), it is very obvious that the ductility can be enhanced using a medium manganese TRIP steel. Therefore, increased elongation with similar levels of ultra high strength of a medium manganese steel in HF condition may provide an alternative to the 22MnB5 grade for some applications. The strength levels of the medium manganese steel after hot forming could likely be enhanced further by using a higher carbon content.

In addition to the mechanical properties, the reheat temperatures used for simulating the HF cycles in the present work are also much lower than in conventional hot stamping due to suppression of Ae3 of the steel by high amounts of manganese. The reheat temperature providing the best mechanical properties (i.e., 700° C.) is at least about 200° C. lower than the temperature used for austenitizing 22MnB5 steel (>900° C.). This would allow for significant energy saving, reducing the costs associated with hot forming, and may have a favorable influence on the behavior of galvanized coating. The invention can also used to hot stamp full- hard starting material to further reduce the number of processing steps. While careful control of the reheat temperatures would be needed in hot stamping of medium manganese steels due to the influence on austenite fraction and stability, the effects of cooling rate should be minimal due to enhanced hardenability resulting from manganese addition. Thus, press quenching may not be a required aspect of microstructure development in these steels.

Finally, the examples provided illustrate the substantial benefits obtained at a 10% manganese level. At much higher levels, the austenite stabilizing effects of manganese can become so great as to minimize the martensitic component of the microstructure, and therefore shift the balance of properties to higher ductility but lower strength.

Ranges have been discussed and used within the forgoing description. One skilled in the art would understand that any sub-range within the stated range would be suitable, as would any number within the broad range, without deviating from the invention.

The foregoing description of the present invention has been presented for purposes of illustration and description. Furthermore, the description is not intended to limit the invention to the form disclosed herein. Consequently, variations and modifications commensurate with the above teachings, and the skill or knowledge of the relevant art, are within the scope of the present invention. The embodiment described hereinabove is further intended to explain the best mode known for practicing the invention and to enable others skilled in the art to utilize the invention in such, or other, embodiments and with various modifications required by the particular applications or uses of the present invention. It is intended that the appended claims be construed to include alternative embodiments to the extent permitted by the prior art.

Claims

1. A method for producing ultra high strength steel alloy, comprising:

providing a ferrous alloy comprising carbon and greater than 2.5 mass % manganese;
annealing the ferrous alloy at a temperature to form an annealed alloy comprising an austenitic microstructure; and
cooling the steel comprising the austenitic microstructure to room temperature to form an ultra high strength steel comprising the austenitic microstructure, and at least one microstructure of ferrite or martensite.

2. The method of claim 1, wherein the annealing the ferrous alloys is performed at a temperature of between about 600° C. and about 800° C.

3. The method of claim 2, wherein the ultra high strength steel comprises about 5 vol. % to about 65 vol. % austenite, between about 35 vol. % and about 90 vol. % ferrite, and the balance martensite.

4. The method of claim 1, wherein the temperature is between about 750° C. and about 900° C.

5. The method of claim 4, wherein the ultra high strength steel comprises about 5 vol. % to about 65 vol. % austenite, between about 35 vol. % and about 90 vol. % martensite.

6. The method of claim 1, wherein the ferrous alloy comprises between about 3 mass % and about 15 mass % manganese.

7. The method of claim 1, wherein a cooling rate following the hot forming step is substantially less than about 30° C./s.

8. A ductile, ultra high strength steel product, comprising between about 2.5 mass % and 15 mass % manganese and wherein the microstructure comprises about 5 vol. % to about 65 vol. % austenite, about 0 vol. % to about 90 vol. % ferrite, and the balance being martensite.

9. The ultra high strength steel product of claim 8, wherein the steel has a grain size of between about 0.1 micron and about 5 micron.

10. The ultra high strength steel product of claim 8, wherein the dislocation density is between about 109 m−2 and about 1010 m−2.

11. The ultra high strength steel product of claim 8, wherein the steel product has a tensile strength of between about 1000 MPa and about 1700 MPa and has a residual total elongation of between about 10% and about 30%.

12. The ultra high strength steel product of claim 8, wherein the ultra high strength steel product has an x-ray diffraction pattern and wherein the x-ray diffraction pattern shows an average FWHM of all ferrite peaks of between about 0.25° and about 1°.

13. The ultra high steel product of claim 8, further comprising an additive selected from the group consisting of an aluminum, a silicon, a carbon, a titanium, a niobium, a vanadium, and combinations thereof.

14. A method for producing ultra high strength steel, comprising:

providing a cold-rolled medium manganese steel comprising between about 2.5 mass % and about 12 mass % manganese;
hot forming the steel at a temperature between about 600° C. and about 900° C. to form a hot formed steel; and
quenching the hot formed steel to form the ultra high strength steel.

15. The method of claim 14, further comprising annealing the steel prior to the hot forming step.

16. The method of claim 15, wherein an annealing temperature is between about 600° C. and about 900° C.

17. The method of claim 14, wherein the ultra high strength steel comprises at least 10 vol. % austenite.

18. The method of claim 14, wherein a cooling rate following the hot forming step is less than about 30° C./s.

19. The method of claim 14, wherein the temperature is between about 600° C. and about 750° C., and wherein the ultra high strength steel comprises about 5 vol. % to about 65 vol. % austenite, between about 35 vol. % and about 90 vol. % ferrite, and the balance martensite.

20. The method of claim 14, wherein the temperature is between about 750° C. and about 900° C., and wherein the ultra high strength steel comprises about 5 vol. % to about 65 vol. % austenite, and the balance martensite.

Patent History
Publication number: 20160312323
Type: Application
Filed: Apr 22, 2016
Publication Date: Oct 27, 2016
Applicant: Colorado School of Mines (Golden, CO)
Inventors: Radhakanta Rana (Alkmaar), Charles H. Carson (Denver, CO), John Speer (Littleton, CO)
Application Number: 15/136,681
Classifications
International Classification: C21D 9/00 (20060101); C22C 38/04 (20060101); C21D 1/18 (20060101); C22C 38/00 (20060101); C21D 8/00 (20060101); C21D 1/26 (20060101); C22C 38/06 (20060101); C22C 38/02 (20060101);